Additive manufacturing of compositionally graded laser deposited titanium-chromium alloys

Journal Pre-proof Additive Manufacturing of Compositionally Graded Laser Deposited Titanium-Chromium Alloys J. Thomas, J.E. Mogonye, S.A. Mantri, D. Choudhuri, R. Banerjee, T.W. Scharf

PII:

S2214-8604(19)32077-9

DOI:

https://doi.org/10.1016/j.addma.2020.101132

Reference:

ADDMA 101132

To appear in:

Additive Manufacturing

Received Date:

1 November 2019

Revised Date:

31 January 2020

Accepted Date:

13 February 2020

Please cite this article as: Thomas J, Mogonye JE, Mantri SA, Choudhuri D, Banerjee R, Scharf TW, Additive Manufacturing of Compositionally Graded Laser Deposited Titanium-Chromium Alloys, Additive Manufacturing (2020), doi: https://doi.org/10.1016/j.addma.2020.101132

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Additive Manufacturing of Compositionally Graded Laser Deposited Titanium-Chromium Alloys J. Thomas, J.E. Mogonye, S.A. Mantri, D. Choudhuri, R. Banerjee, T.W. Scharf * Department of Materials Science and Engineering and Advanced Materials and Manufacturing Processes Institute (AMMPI), University of North Texas, Denton, TX 76203, USA *

Corresponding Author e-mail: [email protected]

Abstract

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Ti-xCr (9≤x≤28) at% graded alloys fabricated via laser engineered net shaping (LENSTM) are potential candidate materials for graded implants in spinal fixation surgeries due to their tunable Young’s modulus. This is achievable by

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various phase transformations on the laser deposited graded alloys by subjugating them to different β-solutionizing heat treatments. The microstructure of the as-deposited (AD) Ti-xCr alloy was comprised of a metastable β-Ti(Cr) matrix phase along with α-Ti and TiCr2 (an intermetallic Laves phase) at different regions of the graded alloy. In

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comparison, the microstructure of the rapidly air cooled (RAC) Ti-xCr alloy contained ω and TiCr2 precipitate phases, while the slowly furnace cooled (SFC) Ti-xCr alloy has α-Ti and TiCr2 precipitate phases along different regions of

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the graded alloys. An extreme change in micro-hardness values at different regions of the RAC and SFC alloys due to these various phase transformations indicates that the alloy fabricated may be qualified as a potential graded implant material. Additionally, the phase evolution along the compositional gradient of the AD Ti-xCr alloy was correlated to

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theories of reheating in the lower build layers, thermal gradient effect, rapid solidification rates, and enthalpy of mixing between Ti-Cr powders. A detailed explanation to variation in microstructures and mechanical properties observed at compositions of Ti~29at%Cr (a lower build layer) and Ti~28at%Cr (an upper build layer) in the AD Ti-xCr graded

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alloy is also provided.

Key Words: Laser Engineered Net Shaping (LENSTM), β stabilized Titanium alloy, Phase transformation, Omega phase, TEM, micro-hardness.

1. Introduction

Recent advancements in medicinal research have led to suggestions that spinal support implant rods with a varying Young’s modulus may be more preferable than conventionally utilized rigid implants [1–4].

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Since surgeons are confined to work in small spaces during spinal implant surgeries, implant rods with minimal degree of spring back are preferred for loading and unloading purposes [5]. Although speculations on utilizing implant rods with higher Young’s modulus as an alternate solution to reduce the degree of spring back during surgeries has been suggested, implants rods with lower Young’s modulus are still in use as they prevent stress shielding effects on the patient’s body [5]. An appropriate solution to the existing dilemma would be the fabrication of implant rods that possess a varying Young’s modulus. Since β-

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Titanium (Ti) alloys possess various phase transformations, and has been utilized as hip implants, they could potentially be used as spinal implant rods if a varying modulus is achieved across the implant. β stabilized Ti – Chromium (Cr) alloys are also known for possessing various desirable properties such as high temperature resistance, wear durability, oxidation resistance, and bio compatibility [6–11]. Interests in Ti-Cr systems as potential bio-medical materials began in the early 1950’s, when Frost et al. [12] discovered that aged and annealed Ti-8wt%Cr exhibited an unexpected brittleness due to the precipitation of hexagonal ω phase. Interestingly, these alloys also form intermetallic Laves phases at higher Cr concentrations [9,13,14]. Despite being brittle in nature, Laves phases are widely recognized for their

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attractive high temperature structural properties [9,13,15]. Studies on binary Ti-Cr alloys at 30, 40, and, 80 at% Cr suggested that enhanced toughness increments and good mechanical properties may be achieved

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with Laves phase nucleation [13]. Zhao et al. [2] reported that the Youngs modulus of solution treated (ST) Ti-Cr (10, 12, and 14 at% Cr) alloys were 82, 68, and 79 GPa, respectively; while, cold rolled (CR) Ti-Cr (10, 12, and 14 at% Cr) alloys were nearly 92, 85 and 82 GPa, respectively. They concluded that the increase

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in Youngs modulus at low Cr at% in ST and CR Ti-Cr alloys is a consequence of athermal and deformation based ω phase transformation; whereas, a decreased Youngs modulus at high Cr at% is a consequence of

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solid solution strengthening [2]. Tensile strength experiments on ST and CR Ti-Cr (10, 12, and 14 at% Cr) alloys revealed that all the ST and CR Ti-Cr alloys possess tensile strengths >700 MPa and >850 MPa, respectively [2]. Furthermore, springback ratios and cytocompatibility of Ti-12Cr alloys were also

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compared to commonly used implant material (316L SS, Ti-6Al-4V, and TNZT) for validating the applicability of Ti-Cr alloys as potential implant rods [2]. Zhao et al [2] reported that: (1) Ti-12Cr alloys possess a smaller springback per unit load applied than TNZT, and as a consequence Ti-12Cr can bend with

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more ease to required shape during operations, and (2) cytocompatibility evaluations reveal that Ti-12Cr possess a greater density of cultured cell than commonly used implant materials (316L SS, Ti-64, and TNZT). In addition, the presence of either ω or Laves phases resulted in increasing Young’s modulus of the material [5,13,16]; while, increasing volume fraction of β stabilized Ti phase reduce the hardness and hence potentially the modulus of the material. It can therefore be hypothesized that by strategically

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engineering the formation of such ω, Laves, and β Ti stabilized phases in the microstructure of a single implant at site-specific locations will assist in the fabrication of a material with a graded Young’s modulus. Importantly, this approach may also provide a solution, which simultaneously satisfies both the patients and doctors requirements [4,17]. Additively manufactured Ti-6Al-4V alloys has been recently studied for texture analysis, compatibility as a functionally graded material, fracture resistance, and orthopedic applications [18–22]. Implants with graded Young’s modulus fabricated via Laser Engineering Net Shaping (LENSTM) process (a variant of direct laser deposition techniques) [23], allows the variation of the composition and by extension the 2

microstructure along a chosen build direction. LENSTM fabrication technique is a flexible additive manufacturing technique that facilitates the development of three dimensional (3D) parts via a layer-bylayer deposition of molten alloy [24–27]. Layer-by-layer deposition of molten alloys over a build length not only facilitates in fabricating compositionally graded alloys, but also aids in maintaining a reasonable control over part dimensions. Even though additive manufacturing process has reduced the need of taking much consideration for various manufacturing constraints (such as shape complexity, hierarchical complexity, material complexity, and functional complexity), the process still possesses multiple issues. Few issues faced when employing additive manufacturing process for fabrication of parts includes the

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following: (1) available materials, (2) geometric limitations (such as minimum wall thickness), (3) dimensional accuracy, (4) surface roughness, and (5) support design [28–30]. However, with technique progress, these issues can start to be addressed, and consequently the manufacturing constraints of additive

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manufacturing processes (i.e. constraints involved during fabrication of alloys via LENSTM) are not discussed in detail in this study [30]. In this study, we have utilized a LENSTM system to deposit a

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compositionally graded Ti-xCr (9≤x≤28) at% binary alloy, with the goal of altering the microstructure and properties along the build direction. An early study involving LENSTM method of alloy fabrication by

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Banerjee et al. [10] concluded that both the microstructural morphology and phase evolution are heavily dependent on solidification time and enthalpy of mixing of pure elemental powders. Zhang et al. [14] used a similar approach to deposit an alloy varying from pure Ti to Ti-60at% Cr, and concluded that although

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formation of TiCr2 intermetallic is a sluggish transformation, rapid solidification during deposition leads to the formation of TiCr2 intermetallic in metastable β-Ti(Cr) at higher Cr content. In other words, the LENSTM fabricated Ti-Cr microstructure of the finished products is expected to undergo several phase

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transformations sequences that can lead to the formation of both metastable and stable phases normally seen in β-stabilized Ti alloys.

Nag et al. [31] and Collins et al. [26] initially studied LENSTM deposited functionally graded β-Ti alloys such as Ti-xTa (0
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of changing chemical composition, (2) effects of heat treatments on the microstructure and micro-hardness as a function of changing chemical composition, and (3) potential use of functionally graded β-Ti alloys as an implant material [26,31]. Depending on the Cr content, Ti-Cr forms both athermal and isothermal ω phase [5]. For example, Liu et al. [5] studied the formation of athermal and isothermal ω phase nucleation in Ti-9Cr-0.2O and concluded that an increase in hardness of the alloy occurs when ω phase nucleates in the β-Ti matrix. Other studies also indicate that Ti-Cr alloys exhibit phase separation prior to the formation of either ω or stable α-Ti phases [10,32]. Narayanan et al. [32] suggested that such phase separations occurs due to the presence of a large miscibility gap, which can clearly be seen in the Ti-Cr phase diagram [32]. 3

Despite such thorough studies on Ti-Cr alloy systems, there is lack of understanding on how the abovementioned transformation pathways will manifest themselves in a compositionally graded laser deposited component. To achieve this, we have deposited a compositionally graded Ti-xCr (9≤x≤28) at% alloy using a LENSTM system. The graded alloy was then sectioned and subjected to various heat treatments followed by subsequent microstructural analysis. Furthermore, they were systematically assessed as a function of increasing Cr content, via electron microscopy observations at multiple length scales coupled with correlative micro-hardness measurements. Additionally, the current study is also focused on establishing a correlation between the microstructural phase evolution and micro-hardness in the as-deposited (AD)

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graded Ti-xCr alloy. Since the microstructural phase evolution in the AD graded Ti-xCr alloy is heavily dependent on various factors (such as rapid solidification rates, melting and mixing of Ti-Cr powders at

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melt pools during fabrication, laser scan speed, reheating effects on lower build layers, etc.), they are expected to be different from what is traditionally observed in as-casted Ti-Cr alloys [6,7,33]. The current study on fabrication and characterization of LENSTM deposited graded Ti-xCr (9≤x≤28) at% alloy (both

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AD and heat treated) not only facilitates in understanding the correlation between microstructure and microhardness at different Ti-Cr compositions, but also assists in improving the methodology of building intricate

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parts rapidly. 2. Material and methods 2.1 Material fabrication

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Ti and Cr powders of 99.9% purity and mesh size of 150 were procured from Alfa Aesar for deposition of the graded alloy. Two containers comprising of 300g of Ti-9at%Cr and Ti-28at%Cr were mixed individually and subsequently rolled on a twin roller ball mill for approximately 4 hours. The rolling

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operation was done to ensure homogeneous mixing between the Ti and Cr powders in each container. The local composition of the graded alloys at different regions are presented in at% throughout the current study. The LENSTM system used for deposition of the graded Ti-xCr (9≤x≤28) at% alloy was an Optomec 750 equipped with a Nd:YAG laser with a wave length of 1.064 µm. A 3D computer aided design (CAD) file of a cylinder consisting of 100 layers was designed and utilized

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to produce a cylindrical alloy with a cumulative height of 25.4 mm and radius 4.8 mm via laser deposition. The height between each layer is designated to be ideally 0.254 mm and contains multiple parallel line hatches. The power of the Nd:YAG laser used for the deposition was set to 480W as a consequence of similar energy being selected by Samimi et al. [34] who also fabricated various LENSTM deposited Ti alloys. The deposition was carried out in an inert Ar atmosphere inside an enclosed glove box with the oxygen level in the glove box maintained below 5 ppm to minimize oxidation during rapid solidification. A Ti-6Al4V (wt%) substrate was used for laser deposition since it allows good adherence with the deposited alloy. A dual powder feeders/hoppers system that deliver powder during deposition were each loaded with Ti4

9at%Cr and Ti-28at%Cr. To achieve a graded alloy deposition, the nominal flow rates for the powder feeders were programmed to change after every 20 layers or 5.08 mm to attain a (3-4) at% increase in Cr compositions. Changing the compositions after every 5.08 mm facilitated in the eventual fabrication of a Ti-xCr (9≤x≤28) at% graded alloy cylinder that is 25.4 mm in height. The programmed powder feed/hopper flow rates, layers covered, and, expected Ti-Cr compositions at the different layers are detailed in Table. 1. Table. 1: Powder feed/hopper flow rates for fabrication of the AD graded Ti-xCr (9≤x≤28) alloy.

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Hopper 1 Hopper 2 Cr Ti Layers Build height flow rate [Ti-9Cr] flow rate [Ti-28Cr] (at %) (at %) (mm) (g/min) (g/min) 0.5 0 9 91 0-20 0-5.08 0.375 0.125 14 86 20-40 5.08-10.16 0.25 0.25 19 81 40-60 10.16-15.24 0.125 0.375 23 77 60-80 15.24-20.32 0 0.5 28 72 80-100 20.32-25.4 Although the AD graded Ti-xCr (9≤x≤28) alloy was ideally designed to possess a composition of

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#

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Ti~9at%Cr at the bottom of the build with a sequential increase in Cr at% after every 5.08 mm, inhomogeneities in the quantity of powder blown from the powder feeds/hoppers lead to acceptable discrepancy in composition at different layers. For example, near the top of the build (~23.5 mm from the

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substrate surface), a local composition of Ti~28at%Cr was observed; while, at a lower region of the build (~20.5 mm from the substrate surface), a local composition of Ti~29at%Cr was observed. Similarly, the

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first 10 layers of deposition was observed to possess a composition of Ti~7at%Cr rather than the programmed Ti~9at%Cr. Even though minor discrepancy in local composition exists at certain layers of the graded Ti-xCr (9≤x≤28) alloy, important information (such as phase transformations at different

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chemical compositions, effects of reheating on the microstructure of the lower build layers, and changes in local micro-hardness as a function of composition and microstructure) at different regions of the graded alloy was revealed from the current study.

Following LENSTM deposition, a Mitsubishi FX 10 wire electrical discharge machine (EDM) was used to separate the AD alloy from the Ti-6Al-4V substrate. The AD alloy was then sectioned vertically along the build length of the alloy to preserve the compositional gradient of the sectioned alloy pieces. The flat

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surface (along the length of the alloy) of one of the sectioned alloy pieces was polished using a progression of SiC abrasive papers from grades 120 to 1200, and then polished to a mirror finish using a colloidal suspension of 0.04 μm silica. Furthermore, the alloy piece was polished with SiC abrasive paper of lower grades for an extended period to prevent any influence of the EDM machining on the microstructure of the alloy. Post mirror polishing, the graded alloy piece was cleaned in an ultrasonic bath of deionized water and ethanol, and subsequently dried using nitrogen gas. Etching was not performed on the surface of the alloy piece as a consequence of the alloy piece being additively manufactured and necessity to determine a

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correlation between microstructure and mechanical property of the AD graded Ti-xCr (9≤x≤28) at% alloy without any bias/influence from external factors. 2.2 Alloy Heat Treatments The second AD graded Ti-xCr (9≤x≤28) at% alloy piece was further sectioned into two more pieces via EDM machining. The sectioning was performed vertically to preserve the compositional gradient along the build length of the two newly sectioned alloy pieces. Each of the newly sectioned alloy pieces were then subjected to β-solutionizing heat treatments at 1000°C for 30 minutes in an inert Ar atmosphere. βsolutionizing at 1000°C was done so as to attain a single β-Ti(Cr) phase throughout the build and relieve

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retained internal stresses developed during the laser depositing process while maintaining the compositional gradient. Following the β-solutionizing treatments, one of the alloy pieces was rapidly air cooled (RAC) by

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placing it in open air; whereas, the other piece was subjected to a slow furnace cooling (SFC) by turning off the furnace power. The alloys were cooled at different rates to determine: (1) the effects of kinetics of phase transformations on the microstructure and mechanical properties exhibited by the graded alloys, and

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(2) plausibility of producing ω phase at low Cr at% via rapid air cooling (usual methods used for production of ω phase in Ti alloys involve quenching and mechanical deformation). Here after, the laser deposited

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specimen is referred as AD, and the homogenized alloys are referred to as RAC or SFC on basis of heat treatment and rate of cooling provided. The polishing procedure detailed in section 2.1 was then used for preparation of the RAC and SFC alloys for various microstructural and mechanical property experiments.

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2.3 Microstructural Characterization

Characterization techniques implemented on the AD and β-solutionized alloys include metallography, elemental analysis, crystal structure analysis, and correlative micro-hardness analysis. A Nova NanoSEM

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230 with energy dispersive spectroscopy (EDS) capability was utilized to attain metallographic information and nominal local area composition at various regions of the alloys. Crystal structure information of the nucleated ω, α-Ti, and TiCr2 intermetallic phase in β-Ti(Cr) matrix were determined by X-ray diffraction (XRD), electron back scatter diffraction (EBSD), and transmission electron microscope (TEM) analysis on TEM lamellas operated in Tecnai G2 F20 at 200kV. The SEM-BSE, EDS, EBSD, and XRD investigations

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were performed along the compositionally graded build length of the AD, RAC, and SFC alloys, respectively. The investigations were performed in such a manner to determine the correlation between microstructure and mechanical properties as a function of changing chemical composition. In order to assess the micro-hardness variation as a function of the programmed compositional and

resultant microstructural variation, a triplicate series of 50 indents of 500 µm spacing were recorded along vertical build direction. Each of the test alloys (AD, SFC, and RAC) were subjected to indentation tests through the length of the graded alloy by utilizing a load of 500g held for 10 seconds at 0.5HV. A triplicate

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series of measurements also assisted in providing a cohesive and meaningful comparison between the changes in micro-hardness trends observed for the test alloys.

3. Results 3.1 X-ray Diffraction 3.1.1

AD alloy

XRD scans of the AD alloy at different compositions along the graded length of the alloy are presented in Fig. 1 (a-c). As shown in Fig.1 (a), at a low Cr at% of Ti~7Cr; α-Ti phase (hexagonal closed packed

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(HCP): P63/mmc space group) is present in a β-Ti(Cr) matrix phase (body-centered cubic (BCC): Im3̅m space group). The phases observed at low Cr compositions in this study were also observed by Ho et al. [6], and Hattori et al. [33]. XRD scans at regions with higher Cr at% of Ti~19Cr (see Fig. 1 (b)) reveal the

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presence of intermetallic TiCr2 (C15 lave phase-face centered cubic (FCC): Fd3̅m space group), α-Ti, and β-Ti(Cr) phases. In contrast, Ho et al. [6] and Hattori et al. [33] observed a retained metastable β-Ti(Cr)

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phase at similar compositions. The upper most layer near the top of the build (possessing a local composition of Ti~28Cr) was observed to have a retained metastable β-Ti(Cr) phase (see Fig. 1 (c), Fig. 3

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(g) and Fig. 4 (d)); while, a few layers beneath the upper layer (possessing a local composition of Ti~29Cr) is observed to have TiCr2 intermetallic phase in a β-Ti(Cr) matrix phase (see Fig. 3 (f) and Fig. 4 (a)). The BSE micrographs of the AD alloy presented in Fig. 3 (a-g) corroborates the XRD scans presented in Fig.1

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the AD graded Ti-xCr alloy.

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(a-c). Table. 2 details the various phases and their crystal structures observed at different compositions of

Figure. 1:(a-c) X-ray diffraction along the compositional gradient of the AD alloy.

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Table. 2: Summary of phases and crystal structures observed along the compositional gradient of the AD alloy via XRD.

Phase

Crystal Structure

Condition

Ti~7Cr

α+β

HCP+BCC

AD

Ti~19Cr

α+β+TiCr2

HCP+BCC+FCC

AD

Ti~28Cr

β

BCC

AD

RAC and SFC alloys

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3.1.2

Alloy (at%)

XRD scans of the RAC alloy at different compositions along the graded length of the alloy are presented in Fig. 2 (a-c). As seen in Fig. 2 (a), definite presence of β-Ti(Cr) and a few unknown broad peaks, later

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identified using TEM and indexed as ω peaks are seen at regions with a low Cr contents of Ti~8Cr. The ω phase indexing in Fig. 2 (a) was also confirmed by other XRD studies of ω phase transformation in β-Ti

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alloys [17,35–39]. As presented in Fig. 2 (b, c), at Ti~11Cr a completely retained β-Ti(Cr) matrix phase was observed; whereas, at Ti~27Cr there was the presence of intermetallic TiCr2 in β-Ti(Cr) matrix phase.

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The morphologies of the phases observed by XRD lineouts are also presented in the micrographs from Fig. 5 (a-c), and Fig. 6 (a). These findings corroborate conclusions drawn by Ho et al. [6] who assessed the microstructure of cast Ti-Cr alloys and determined that cast Ti-Cr alloys with Cr content less than 10 wt%

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Cr possess a dual ω+β phase; whereas, those with Cr content between 25 and 30 wt% show a retained single β-Ti phase [6]. Qualitative assessment of Fig. 5 (c, f) suggest that both RAC and SFC graded Ti-Cr alloys show the presence of nucleated C15 FCC structured TiCr2 intermetallic Laves phase in β-Ti(Cr) matrix at

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increased Cr at%.

Figure. 2: X-ray diffraction along the compositional gradient of (a-c) RAC alloys, and (d-f) SFC alloys.

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XRD scans of SFC alloys presented in Fig. 2 (d-f), indicate the presence of a dual α+β phases for Ti~7Cr and Ti~11Cr. These observations are in accordance to the observations made by Ho et al. [6] and Hattori et al. [33]. From the lineout presented in Fig. 2 (f), and corresponding micrograph in Fig. 5 (f) it is evident that intermetallic C15 TiCr2 Laves phases surrounded by α-Ti are embedded in a β-Ti(Cr) matrix phase at a nominal composition of Ti~27Cr at%. The observance of an intermetallic surrounded by α-Ti in β-Ti(Cr) matrix phase is in contradiction to the observation of a fully retained β-Ti(Cr) phase seen in casted Ti-Cr alloys by Ho et al. [6] and Hattori et al. [33]. Hence from these XRD experiments, it can be concluded that depending on the alloy composition analyzed, SFC alloys show presence of α-Ti, β-Ti(Cr) matrix

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phase, and C15 TiCr2 intermetallic Laves phase. The metallographic images presented in Fig. 5 (a-f) at various compositional regions corroborates the XRD lineouts presented in Fig. 2 (a-f). The various phases observed at different local compositions for both RAC and SFC alloys are summarized in Table. 3.

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Table. 3: Summary of phases observed along the compositional gradient of RAC and SFC alloys via XRD.

Phase(s)

Crystal Structure

Condition

Ti~8Cr

ω+β

Hexagonal+BCC

RAC

Ti~11Cr

β

BCC

RAC

Ti~27Cr

β+TiCr2

BCC+FCC

RAC

Ti~7Cr

α+β

Ti~11Cr

α+β

Ti~27Cr

α+β+TiCr2

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SFC

HCP+BCC

SFC

HCP+BCC+FCC

SFC

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3.2.1 AD alloy

HCP+BCC

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3.2 Microscopic Characterization

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Alloy (at%)

Fig. 3 (a-g) represent the back-scatter electron (BSE) micrographs at site specific locations of the AD alloy. Fig. 3 (a-c) represents micrographs at low Cr compositions. At these regions, β-Ti(Cr) is the matrix phase and α-Ti is the precipitate phase. From Fig. 3 (a-c), it is evident that an increase in Cr at% along the length of the graded alloy decreases the nucleation of the α-Ti precipitate phase. Visual assessments of Fig. 3 (a-c) reveal the quantity of α-Ti in the β-Ti(Cr) decreases with increasing Cr content. At around Ti~19Cr,

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C15 TiCr2 intermetallic phase begins to nucleate in the β-Ti(Cr) matrix phase along with negligible amounts of α-Ti phase (see Fig. 3 (d)). Fig. 3 (e), clearly shows the nucleation of TiCr2 intermetallic phase along the grain boundary of the β-Ti(Cr) phase when the local composition increases to Ti~23Cr. Fig. 3 (f) reveal that at Ti~29Cr (a lower build layer) of the AD alloy located ~20.5 mm from the bottom of the substrate surface has an extensive intermetallic TiCr2 phase precipitation within a β-Ti(Cr) matrix phase. In contrast, Fig. 3 (g) reveal that Ti~28Cr (an upper build layer) of the AD alloy located ~23.5 mm from the bottom of the substrate surface has a retained metastable β-Ti(Cr) phase.

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Figure. 3: (a-g) BSE micrographs along compositional gradient of the AD alloy.

Fig. 4 (a-f) presents EBSD maps of the AD alloy at local compositions of Ti~29Cr (a lower build layer) and Ti~28Cr (an upper build layer). The EBSD scans reveal that the lower build layer (Ti~29Cr) possess a C15 TiCr2 phase in a β-Ti matrix phase, while the upper build layer (Ti~28Cr) possess a retained metastable β-Ti matrix phase. Fig. 4 (a-c), indicate that the C15 TiCr2 intermetallic phase is randomly oriented and

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occupies a phase fraction of 0.351 of the area analyzed. The reheating of the lower build layers is believed to have facilitated the nucleation of the TiCr2 intermetallic phase in the β-Ti matrix phase at Ti~29Cr (see Fig. 4 (a-c)). Fig. 4 (d-f) representing EBSD maps at Ti~28Cr (upper build layer) show a retained metastable β-Ti phase with partial texturing. Being the upper most region or final build layers of the AD alloy, the region is devoid of effective cycles of reheating post rapid solidification. Since the upper build layer (Ti~28Cr) is devoid of reheating cycles, they have a retained metastable β-Ti matrix phase. Fig. 4 (f), reveals the partial texturing of the β-Ti phase along {001} pole and partial clustering of grains in the [001]

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direction. Similar effects of partial texturing in LENTM deposited Ti-Cr alloys was also observed by

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Banerjee et al. [10].

Figure. 4: EBSD mapping of the AD alloy at Ti~29Cr (lower build layer) representing (a) Phase fraction Map, (b) TiCr 2 inverse pole figure (IPF), (c) TiCr2 Texture plot; and Ti~28Cr (upper build layer) representing (d) SE image, (e) β -Ti IPF, (f)

3.2.2 RAC and SFC alloys

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β -Ti Texture plot.

Fig. 5 (a-c) presents the BSE micrographs of three site specific regions on the graded RAC alloy. Fig.

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5 (a) presents a micrograph at a low Cr at%. At these regions, β-Ti(Cr) is the dominant matrix with embedded ω phase. Although the small precipitate size of the athermal ω phases (~5 nm) render them unviewable by BSE imaging, dark field TEM (DF-TEM) images presented in Fig. 6 (a) clearly shows the presence of these precipitates. Fig. 5 (b) represents a completely retained β-Ti(Cr) phase at Ti~11Cr, while Fig. 5 (c) indicates extensive nucleation and growth of TiCr2 in the β-Ti(Cr) matrix at Ti~29Cr. The BSE

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micrographs of various compositions coupled with EDS analyzed local chemistry for SFC alloys are presented in Fig. 5 (d-f). At low Cr at% concentration of Ti~7Cr, there is the presence of large scale homogeneous α-Ti(Cr) phase as dendritic structures within a β-Ti(Cr) matrix. Further visual assessments of Fig. 5 (d-f) suggests that the quantity of α-Ti decreases with respect to increasing Cr at% along the length of the graded alloy. At Ti~27Cr, there is extensive TiCr2 intermetallic Laves phase along with α-Ti phase at its phase boundaries.

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Figure. 5: BSE micrographs along the compositional gradient of (a-c) RAC alloys, and (d-f) SFC alloys.

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Selected area electron diffraction (SAED) patterns and DF-TEM images of the RAC specimens produced at a composition of Ti~9Cr indicate the presence of well-structured athermal ω phase within a β-

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Ti(Cr) matrix. Fig. 6 (a, b) represent the respective scattered diffraction of the athermal ω phase as white spots in DF-TEM image mode, and SAED pattern of athermal ω phase indexed inside the β-Ti(Cr) matrix phase. The SAED patterns of the athermal ω phase was acquired along the [011] zone axis of β-Ti(Cr). Fig.

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6 (c-f) represents micro-diffraction patterns from the three zone axes of TiCr2, and bright field TEM (BF-

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TEM) micrographs from SFC alloy specimens obtained at a local composition of Ti~21Cr at%.

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Figure. 6: (a, b) DF-TEM micrograph and corresponding SAED pattern of athermal ω phase indexed in [011] zone axis of βTi(Cr) matrix phase of RAC alloy specimen at a nominal composition of Ti~9Cr, (c-f) BF-TEM image and corresponding micro-diffraction patterns of TiCr2 phase indexed at various zone axes of the SFC alloy specimen at Ti~21Cr. The micro-

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diffraction patterns reveal that the TiCr2 has a FCC crystal structure. The red dotted box encircling the precipitate in fig. 6 (c) indicates the region from where the micro-diffraction patterns were procured.

The electron micro-diffraction patterns of TiCr2 phase presented in Fig. 6 (d-f) show typical FCC

arrangement of diffraction spots in a reciprocal lattice at [111], [011] and [112] zone axes. The FCC structure was further validated using XRD observations to determine that the TiCr2 intermetallic is a C15 Laves phase with Fd3̅m symmetry. Fig. 7 (a) represents a high angle annular dark field (HAADF) micrograph produced in scanning TEM (STEM) mode. As presented in Fig. 7 (a, b), a high resolution EDS line scan is also carried out along the β-Ti(Cr) → α-Ti → TiCr2 phases in order to determine the changes in element composition (at%) along the three phases of the SFC alloy. The β-Ti(Cr) phase appeared to be 13

a mix of Ti and Cr with a Ti to Cr ratio of 80:20 at%; whereas, the α-Ti phase was identified to be near 100% Ti. The TiCr2 intermetallic being predominantly occupied by Cr with small amounts of Ti appeared

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to be in a ratio of 60:40 at% and not the stoichiometric 2:1 ratio of 67:33 at%.

Figure. 7: (a) STEM-HAADF micrograph of SFC alloy with STEM-EDS line scan on micrograph passing through matrix βTi(Cr), a-Ti, and TiCr2 phases, and (b) Compositional profile of STEM-EDS line scan presented in Fig. 7 (a).

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Furthermore, EBSD mapping of the SFC alloy at a local composition of Ti~22Cr presented in Fig. 8 (a-d) validates the presence of the three phases and revealed a probable Burgers vector relationship between β-Ti and α-Ti phases. The average confidence index of the cleaned micrograph is 0.69, with 0.401, 0.248,

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and 0.350 fraction of the micrograph being occupied by β-Ti(Cr), α-Ti and TiCr2 phases, respectively. Furthermore, a closer observation of the EBSD pole figures (PF) presented in Fig. 8 (c, d) indicate that certain spots in the basal {0001} α-Ti pole (marked by red circle’s) aligns with certain spots in the {101} β-Ti(Cr) pole (marked by red circle’s). This suggests that there exists a potential Burgers vector relationship between the two phases. Since clustering of α-Ti precipitates around TiCr2 precipitates in β-Ti(Cr) matrix was observed from both EBSD and STEM micrographs of the SFC alloys, there is a possibility that at

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higher Cr at%, the TiCr2 intermetallic phase nucleation is followed by the subsequent formation of α-Ti phase, which is a Cr depleted phase. In addition, due to the α-Ti laths forming after the nucleation and growth of TiCr2 intermetallic, it can be assumed that the α-Ti phase may have potentially selfaccommodated with respect to the β-Ti(Cr) matrix in order to reduce the overall straining in the alloy system [40–42]. Self-accommodation of α-Ti laths to minimize the total transformation shape strain and strain energy by clustering in specific crystallographic orientations was also observed by Nag et al. [40]. Nag et al. [40] provided a detailed reasoning behind the clustering and orientation of various individual α-Ti 14

variants in β matrix phase of Ti5553 alloys. They observed that three α-Ti variants from a furnace cooled Ti5553 alloy grew with {0001}α//{011}β relationship with each variant lying at an angle of 60˚ with respect to each other, and sharing a common <11-20> α //<111> β direction [40]. Correlating the observations seen via EBSD in this study with respect to observations made by Nag et al. [40], it can be assumed that the orientation relation of the α-Ti with respect to β-Ti matrix in the SFC alloys is {0001}α//{101}β. However, considering that only a small area of grains was analyzed, a conclusive and cohesive assessment on the orientation relationship between the phases in the entire graded SFC alloy cannot be provided. Further

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detailed assessment on orientation relationship between phases is beyond the scope of this study.

Figure. 8: EBSD mapping of SFC alloy at Ti~22Cr showing (a) Phase fraction map, (b) IPF of α-Ti phase, (c) PF of α-Ti phase, (d) PF of β-Ti phase.

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3.3 Micro-hardness Evaluation 3.3.1 AD alloys Vickers micro-hardness indentation tests at various locations of the AD alloy is presented in Fig. 9 (a, b). Fig. 9 (a), reveals the change in micro-hardness as function of increasing Cr at% (in blue); and, change in micro-hardness as a function of alloy build length (in red). Furthermore, the error associated to hardness value along the build length is presented as a shaded red region along the micro-hardness vs build length trend line (in red). The micro-hardness tests revealed that the presence of the TiCr2 intermetallic phase at

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regions with high Cr at% caused a spike in the average micro-hardness to ~450HV.

Figure. 9: (a) Change in micro-hardness as a function of increasing Cr content (in blue) and build length (in red) in the AD

alloy, and corresponding BSE micrographs of Vickers indents at different compositions of the AD alloy (see markers labelled 1-6), (b) low magnification BSE micrographs near Ti~22Cr showing the Vickers indents on the TiCr2 intermetallic phase and the β -Ti(Cr) matrix phase. The markers labeled (1-6) in the graph in Fig.9 (a) correspond to the makers labelled (1-6) in the

BSE micrographs in Fig. 9 (a, b). These markers represent the approximate Cr at% and micro-hardness of the Vickers indents presented in the BSE micrographs.

As seen from Fig. 9 (a), regions with low Cr at% (Ti~(7-9)Cr and Ti~(12-19)Cr), show an average micro-hardness variation between ~(350-360)HV. At Ti~22Cr a large deviation in micro-hardness is 16

observed due to indents measuring hardness at grain boundaries where intermetallic TiCr2 phases are present, and within single phased β-Ti(Cr) stabilized grains (see regions marked 4 in Fig. 9 (a, b)). The indentation measurements at these regions also indicate that the presence of intermetallic increase the micro-hardness; while, the presence of retained metastable β-Ti(Cr) grains decrease the micro-hardness. The BSE micrographs of indents presented in Fig. 9 (a, b) in conjunction with the micro-hardness measurement at various compositions corroborate the statement on increased and decreased micro-hardness due to the presence of TiCr2 and β-Ti(Cr) phases respectively. Furthermore, the micro-hardness vs Cr at% plot presented in Fig. 9 (a) also indicates that at around Ti~29Cr, there is an increase and decrease in micro-

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hardness. The increased hardness is the due to the nucleation of TiCr 2 intermetallic phases at the lower build layers; while, the decreased hardness due to the presence of a retained metastable β-Ti(Cr) phase at

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the upper build layers respectively. The correlated variation in microstructure and micro-hardness at the lower- and upper-build layers is clearly seen in Fig. 9 (a) (see markers labelled 5 and 6). 3.3.2 RAC and SFC alloys

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The various microstructural phases evolved in the graded RAC alloys indicate there are likely changes in the hardness along the length of the graded alloy. Fig. 10 (a) represent the local hardness in RAC alloy

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as a function of increasing Cr at% (blue scatter plot). From Fig. 10 (a), it can be observed that the highest hardness in the graded RAC alloy is at Ti~9Cr. The presence of the harder athermal ω phase causes a spike in hardness and results in a local area with hardness ~550HV. Although the athermal ω phase nucleation

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accounted for a substantial increase in ~300HV, a decrease in hardness is observed when the athermal ω phase is depleted by increasing Cr at% [6,7,38]. At higher Cr at% of Ti~(12-18)Cr, 100% β-Ti(Cr) phase retention observed in both BSE micrographs and XRD lineouts has resulted in a decrease in hardness to

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~360HV [7]. At Ti~(23-28)Cr an increase to around ~500 HV is observed owing to the nucleation of TiCr2 intermetallic phase at higher Cr at%. For RAC alloys, the presence of α-Ti phase nucleation along with TiCr2 intermetallic phase in β-Ti(Cr) is not observed despite β-Ti(Cr) phase decomposition being a eutectoid reaction. The micrographs of the Vickers indents presented in Fig. 10 (b), illustrates that the highest hardness in RAC alloys is at low Cr at% where athermal ω phase is present. With higher hardness

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in these regions, there is corresponding brittleness indicated by presence of shear bands surrounding the indents, a sign of strain localization. These bands (marked with white arrows) are clearly seen in the micrograph labeled Ti~8Cr in Fig.10 (b). The other images of the indents at different regions of the RAC alloys do not show this deformation as there is no athermal ω phase nucleation at increased Cr at%.

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Figure. 10: (a) Change in micro-hardness as a function of increasing Cr content in RAC and SFC alloys, (b) BSE images of

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micro-hardness indents at different compositions of RAC alloy, (c) BSE images of micro-hardness indents at different compositions of SFC alloy.

The various microstructural phases evolving in SFC alloys also imply there are changes in micro-

hardness as a function of Cr composition. Fig. 10 (a) shows the hardness as a function of Cr composition for the SFC alloys (in red). The hardness of the SFC alloy at low Cr at% is observed to be ~400HV due to homogeneous α-Ti phase nucleation. The decrease in α-Ti phase nucleation as a response to increase in Cr at% results in a decrease in hardness to ~360HV at Ti~(14-18)Cr. The sudden increase in hardness at Ti~(23-28)Cr to ~500 HV is likely due to the presence of three phase regions (α-Ti, β-Ti(Cr) and TiCr2). 18

Fig. 10 (c), illustrates that the size of the Vickers hardness indentation decreases with increasing Cr at%. In addition, while hardness increases with increasing Cr content due to TiCr2 phase formation, there is no evidence of shear bands propagating along corners of the indents seen in Fig. 10 (b). Furthermore, this indicates that the SFC alloy maintains an increase in hardness without much of a compromise in toughness. 4. Discussion The AD graded Ti-xCr alloy shows precipitation of α-Ti phase in a β-Ti(Cr) matrix phase at low Cr at%, and C15 TiCr2 intermetallic phase in a β-Ti(Cr) matrix at high Cr at% respectively (see Fig. 1 (a-c) and Fig. 3 (a-g)). As per the Ti-Cr binary phase diagram, Ti~19Cr and Ti~23Cr are hyper eutectoid

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compositions, and are therefore expected to have a large quantity of TiCr2 intermetallic phase and a reduced quantity of α-Ti nucleate in a metastable β-Ti(Cr) matrix phase. Early Ti-Cr studies indicate that the

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eutectoid reaction in Ti-Cr alloys is sluggish, and the sluggish nature hinders the nucleation of the TiCr2 intermetallic phase at Cr concentrations below 5 at% [43–46]. Fig. 4 (d-f) representing EBSD maps at Ti~28Cr (an upper build layer), revealed partial texturing of the β-Ti phase. A similar effect of partial

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texturing in LENSTM deposited alloys was also reported by Banerjee et al. [10]. However, considering that only a small area of grains was analyzed, a cohesive conclusion on the partial texturing of the β-Ti in the

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entire AD alloy cannot be provided. A more detailed texture analysis is beyond the scope of this study. The quantity of α-Ti and TiCr2 phases observed at different regions of the AD graded Ti-Cr alloys is believed to be dependent on the flow of heat from the melt pool to the substrate. Since the heat produced at

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the melt pool during alloy deposition flows from the upper- to lower-layers of the build, the solidified lower build layers are subjected to cycles of reheating [47–49]. As a consequence, the reheating cycles affect the final microstructure and micro-hardness regions of the AD part build. Kelly et al. [50,51] studied the effects

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of reheating cycles on the microstructural properties of deposited Ti-6Al-4V. Their attempt to quantify the process of nucleation, growth, and transformation of new phases from the top to bottom of a Ti-6Al-4V specimen was well explained via experimental analyses and correlated numerical modelling [50,51]. Several thermo-kinetic and thermo-analytic models on laser deposited alloys has also been developed to describe the effect of reheating cycles during alloy deposition in a greater detail [47,49,52–55]. Most of the

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current model’s on laser deposited materials revealed that the heat generated at the melt pool during laser deposition migrates from the upper- to lower- layers of the specimen and is then dissipated into the substrate [47,50,55]. In a similar context, form this study we observe that a lower build layer (~20.5 mm from the substrate surface) with a local composition of Ti~29Cr is subjected to reheating via to flow of heat from the upper build layers. As seen in Fig. 4 (a-c), reheating causes the β-Ti(Cr) matrix phase in the lower build layers (Ti~29Cr) to decompose into β-Ti(Cr) + TiCr2 intermetallic phase; while, the upper build layer (~23.5 mm from the substrate surface) with a local composition of Ti~28Cr not being subjected to any reheating possess a fully retained metastable β-Ti(Cr) phase (see Fig. 4 (d-f)). 19

Another perspective on phase transformations in laser deposited Ti-Cr alloys was provided by Banerjee et al. [10] who focused on the enthalpy of mixing between Ti-Cr powders. Studies show that Ti-Cr systems possess a negative enthalpy of mixing (-3000 cal/g atom), with the negative sign indicating that heat is released during the mixing process [10,56]. Banerjee et al. [10] and Hofmeister et al. [57] suggested that during the LENSTM deposition process, the temperature at the melt pool can increase above the usual melt temperature due to a localized heating effect from the enthalpy of mixing between the Ti-Cr powders [10,57]. They also stated that if the rate of solidification is assumed to be proportional to the temperature difference of the surrounding substrate and melt pool, the higher temperature at the melt pool would

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facilitate a faster solidification rate [10]. As a consequence, laser deposited Ti-Cr alloys (especially at higher Cr at% content) possess a negative enthalpy of mixing and show a rapidly solidified microstructure post

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deposition [10]. Based on the assessment provided by Banerjee et al. [10], it is hypothesized that regions with a high Cr at% possess a more negative enthalpy of mixing in comparison to those with a low Cr at%. As a consequence, regions with a high Cr at% have higher solidification rates than its counterparts at low

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Cr at%. The variation in solidification rates at different regions of the AD alloy facilitates the high Cr at% regions (with enhanced solidification rates) have a greater yield of C15 TiCr2 intermetallic phase nucleation,

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and low Cr at% regions (with reduced solidification rates) have a greater yield of α-Ti phase nucleation [10]. Similar conclusions were also suggested by Banerjee et al. [10] who studied the phase evolution in laser deposited Ti-Cr alloys. From the current study on phase transformations in the AD graded Ti-xCr

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alloy it can be concluded that (1) the phase transformations at lower build layers in the AD alloy are affected by reheating cycles; while, the upper build layers do not undergo any reheating, and (2) the phases observed at different regions of the AD alloy is dependent on rapid solidification rates and enthalpy of mixing of the

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Ti-Cr powders during the laser deposition process [46,48]. The RAC graded alloy is believed to have followed a diffusionless phase transformation route due to the kinetics involved from rapid air cooling of the alloy from a β-solutionized temperature. The rapid air cooling is believed to have halted the α-Ti phase nucleation, which is predominantly a diffusion governed nucleation. At Ti~(9-10)Cr, a 100% retained metastable β-Ti(Cr) phase was observed from the micrographs

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presented in Fig.5 (a-c); however, XRD, TEM, and micro-hardness analysis, indicate the presence of a wellstructured wide spread athermal ω phase within the β-Ti(Cr) matrix phase at low Cr at%. Although no literature can be found on the composition for initiation of ω phase within the β-Ti(Cr) matrix of Ti-xCr alloys, existence at low Cr at% has been observed in other studies that involved Ti-Cr alloys [5,32,34]. It has also been suggested that Ti-Cr alloy system having a spinodal decomposition of the β phase, would prefer β1 and β2 metastable phases to transform into either β + α or β + ω upon cooling from elevated temperatures [32,43]. Although other studies have noted 100% β phase retention and ω phase nucleation occurring from quick cooling by water quenching [5], it is interesting to note that similar results were 20

achieved in this study when the alloy was rapidly air cooled. Isothermal ω phase is most commonly observed in β-Ti alloys that are subjected to thermal treatments, while athermal ω phase is observed to form in retained β-Ti phase when quenching from elevated temperatures that does not activate martensitic phase transformations [58,59]. Although heterogeneous mechanisms such as the phase separation of the β-Ti phase and formation the ω phases facilitating the heterogeneous nucleation of the α-Ti phase has been reported in other works [58–62], it has not been observed in this study. Furthermore, Narayanan et al. [32], Liu et al. [5], and Samimi et al. [34] reported the formation of ω phase in the metastable β-Ti phase of TiCr alloys that had low Cr content. The phase diagram of Ti-Cr binary alloy system indicates that Ti~12Cr

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is the approximate region where eutectoid reaction of β-Ti(Cr) into α-Ti and TiCr2 phases occur assuming equilibrium conditions are satisfied. Since rapid air cooling of heat treated alloys is technically a non-

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equilibrium based diffusionless process, the suggested phase transformation of β-Ti(Cr) into α-Ti and TiCr2 is believed to be halted. Hence, based on literature available and observations produced from this study, the assumed decomposition pathway for β-Ti(Cr) phase for RAC alloys at regions of higher Cr compositions

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can be assumed to be:

β → coherent β1 + β2 →incoherent β1 + β2→ incoherent β1 + β2 + grain boundary precipitates → stable β

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+ TiCr2 [63].

For composition’s ranging between Ti~(11-19)Cr, a 100% retained β-Ti(Cr) phase was confirmed by visual BSE micrograph assessments and XRD analysis. These observations are in accordance with

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observations made by Welsch et al. [44] and Luhman et al. [64]. In this study, the initial formation of intermetallic TiCr2 was also observed to occur at grain boundaries where the nominal local composition is

Cr content.

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Ti~24Cr, and the volume fraction of TiCr2 phase was also observed to increase as a function of increasing

The SFC alloy is believed to follow a diffusion propagated phase transformation route due to the kinetics involved from slow furnace cooling of the alloy from β-solutionized temperature [65]. The slow cooling enables the α-Ti phase to nucleate and grow throughout the SFC alloy. Although limited, the presence of α-Ti phase is observed at higher Cr concentrations. Cr being a β stabilizing element is believed

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to have suppressed the extensive formation of the α-Ti phase at higher Cr at%. The initial intermetallic TiCr2 formation was observed at Ti~14Cr to occur at grain boundaries and continues to grow with increase in Cr content. The presence of three phases creates a region of high hardness and increases the hardness to ~500 HV. Based on available literature and current observation, the probable decomposition pathway for β-Ti(Cr) phase at high Cr content regions in the SFC alloy is assumed to be : β → β + TiCr2 → α + TiCr2 [63]. Although the formation of TiCr2 at higher Cr content has been suggested to be a sluggish reaction [14,43], the graded SFC alloys showed a dominant presence of TiCr2 phases enclosed by a depleted α-Ti 21

region in a β-Ti(Cr) matrix (see Fig. 5 (f), Fig. 7 (a), and Fig. 8 (a)). TEM and EBSD analysis at high Cr content regimes of the SFC alloys suggest that the TiCr2 intermetallic nucleates and grows within the βTi(Cr) matrix phase followed by the formation of Cr depleted α-Ti phase around the TiCr2 phase. The TiCr2 intermetallic containing a higher at% of Cr than α-Ti nucleates by absorption of Cr content in the matrix alloy. The absorption of Cr from the β-Ti(Cr) matrix creates a Cr depleted and Ti enriched region around the intermetallic boundary. This facilitates the eutectoid reaction to proceed, and form α-Ti phase around the intermetallic phase. Fig. 7 (a, b) representing a STEM micrograph and high-resolution STEM-EDS line scan at Ti~21Cr provide corroborative evidence to the above-mentioned theory.

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From the current study, it is evident that the rate of cooling affects both the microstructure and mechanical properties of the compositionally graded laser deposited alloys. Qiu et al. [66] recently reported

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on the sensitivity of microstructure and mechanical property in β stabilized Ti alloys as a function of various processing parameters during fabrication via laser powder bed fusion (L-PBF). Their assessments revealed the following: (1) small powder layer thickness and low energy density facilitated formation of β columnar

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grains, while an increased energy density facilitated grain growth and texturing, (2) increased powder layer thickness facilitated columnar-to-equiaxed grain transition and reduced texturing, (3) presence of athermal

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ω precipitates in all the as-fabricated specimens, and (4) fine columnar grain structure increased specimen strength and ductility, while hybrid grain structures facilitated intergranular fracturing and reduced ductility. In comparison to the observations made by Qiu et al. [66], the AD alloy assessed in this study

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revealed the following: (1) a micro-hardness of ~360HV at regions with low Cr at% where α-Ti and βTi(Cr) are the stable phases, and an increase in micro-hardness to ~450HV when TiCr2 phase forms as a consequence of increase in Cr at% and cycles of reheating, and (2) effects of reheating cycles on the lower

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build layers (~20.5 mm from the substrate surface) of the alloy leading to nucleation of C15 TiCr2 intermetallic phase with a micro-hardness of ~450HV in comparison to the retained metastable β-Ti(Cr) phase present at the upper build layers (~23.5 mm from substrate surface) of the alloy with micro-hardness of ~400HV. In contrast to the study reported by Qiu et al. [66], the presence of athermal ω phase facilitating an increase hardness was not observed in the AD alloy in this study. Although a direct correlation between

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the assessments provided by Qiu et al. [66] and the current study cannot be made, both studies indicate that the processing parameters, changes in chemical composition, and reheating cycles on lower build layers of laser deposited alloys severely affect the correlated microstructure and mechanical property exhibited by the alloys.

A clear correlation between microstructure and mechanical properties in RAC and SFC alloys is seen at low Cr content regions (see Fig. 10 (a)). The RAC alloy at Ti~(7-9)Cr possessing ω + β-(Ti, Cr) phases had an increased hardness of ~525HV in comparison to SFC alloy that possessed α-Ti + β-(Ti, Cr) phases and a hardness of ~400HV. At a local composition of Ti~(12-18)Cr, both SFC and RAC alloys possess 22

similar hardness of ~360HV despite having varying microstructures. The decrease in hardness for the RAC alloy is due to 100% β-(Ti, Cr) phase retention; whereas, SFC alloys observe a decrease α-Ti phase nucleation and growth. At increased Cr content of Ti~(23-28)Cr, the RAC alloy possessing β-(Ti, Cr) + TiCr2 phases possess a local hardness of ~475HV; whereas, the SFC alloy at similar compositions possessing α-Ti + β-(Ti, Cr) + TiCr2 phases has a local hardness of ~525HV. The increased hardness in the SFC alloy at high Cr content can be attributed to the presence of these three phases. As mentioned previously, increasing hardness often times is indicative of increasing Young’s modulus [1,2]. Thus, a varying Young’s modulus in the SFC and RAC compositionally graded Ti-Cr alloys make them attractive

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candidates for implant rods or other biomedical implants. 5. Conclusions

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The correlation between microstructure and mechanical properties for the AD graded Ti-xCr alloy and Ti-xCr alloys that were homogenized and cooled at differed rates (RAC and SFC alloys) was investigated systematically by coupling XRD, SEM, EDS, EBSD, and TEM characterization along with micro-hardness

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Vickers indentation. The major conclusions drawn from this study are:

1. At low Cr content of Ti~(7-9)Cr at%: The AD alloy show α-Ti phase nucleation in a β-Ti(Cr)

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matrix phase. The micro-hardness for the AD alloy at regions with a low Cr at% has been approximated to ~360HV. In contrast, the RAC alloy possess a ω + β-Ti(Cr) microstructure. Rapid air cooling is believed to have halted the formation of diffusion based α-Ti phase and promoted the

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formation of diffusionless athermal ω phase. The micro-hardness of the RAC alloy at these regions has been approximated to ~525HV. The SFC alloys at similar compositions possess a α-Ti + βTi(Cr) microstructure. The micro-hardness of the SFC alloy at these regions has been approximated

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to ~400HV.

2. At local compositions of Ti~(12-18)Cr at%: Minimal α-Ti phase nucleation in β-Ti(Cr) matrix is observed for the AD alloys with a local micro-hardness approximated to ~350HV. The RAC alloy observe a 100% β-Ti(Cr) phase retention with micro-hardness reverting to ~360HV, while the SFC alloy at similar local compositions of Ti~(12-18)Cr at% showed decreased α-Ti phase nucleation

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in β-Ti(Cr). Being largely β-Ti(Cr) phase, the micro-hardness for the SFC alloy reduces to ~360HV. 3. At Ti~(23-28)Cr at%: The RAC alloy shows the nucleation of TiCr2 phase within β-Ti(Cr) matrix phase. As a consequence, an increase in hardness to ~475HV is observed. The SFC alloys on the other hand possess an increase in hardness to ~525HV. The increased hardness can be attributed to the presence of three phases (TiCr2, α-Ti, and, β-Ti(Cr)) at these regions. In comparison, the AD alloy shows an increase in hardness to ~450HV at similar compositions. This is due to the nucleation of the hard C15 TiCr2 intermetallic phase in the β-Ti(Cr) matrix phase.

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4. Interestingly, the AD alloys, at Ti~29Cr (a lower build layer ~20.5 mm from the substrate surface) has an increased hardness of ~450 HV due to the nucleation of C15 TiCr2 lave phase; while, at Ti~28Cr (an upper build layer ~23.5 mm from substrate surface) shows a reduced hardness of ~400HV due to retention of the metastable β-Ti(Cr) phase with minimal or no nucleation of TiCr2 intermetallic phase. From this study, it is evident that the reheating cycles at the lower build layers, rapid solidification at various layers, and enthalpy of mixing between elemental Ti-Cr powder blends blown from the hoppers during laser deposition are a few of the determining factors that need to be considered when designing and

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fabricating a graded alloy part for potential commercial application. Furthermore, the importance of enthalpy of mixing of element powders at the melt pool and their effects on solidification rates has also

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been discussed in detail. From this study, it is also evident that the different phase transformations appearing in the heat-treated alloys also induce drastic changes in mechanical properties, such as indentation hardness. Hence, depending on the properties required, Ti-Cr graded implants with specific characteristic hardness,

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and likely Young’s modulus, can be engineered via LENSTM. Formation of ω or TiCr2 intermetallic phase regions exhibiting high hardness, and β-Ti(Cr) phase regions exhibiting lower hardness indicate that varied

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mechanical properties can be achieved on a single specimen depending on the site-specific composition and phases tailored during the deposition. Additionally, even though the current study revealed the possibility of tailoring site-specific mechanical properties in a LENSTM fabricated alloy via varying the

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local chemical compositions and heat treatments provided, caution must be taken when assessing the final layers of the laser deposited alloy. The final layers of the LENSTM deposited alloy in comparison to all other layers can display drastically different properties as a consequence of a significantly different thermal

layers.

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history experienced, and hence traditional subtractive machining may be required for removal of the final

Declaration of interests

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The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgment The authors acknowledge the Materials Research Facility (formerly Center for Advanced Research and Technology) at the University of North Texas for access to the material characterization facilities.

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References [1]

M. Nakai, M. Niinomi, X. Zhao, X. Zhao, Self-adjustment of Young’s modulus in biomedical titanium alloys during orthopaedic operation, Mater. Lett. 65 (2011) 688–690. doi:10.1016/j.matlet.2010.11.006.

[2]

X. Zhao, M. Niinomi, M. Nakai, J. Hieda, T. Ishimoto, T. Nakano, Optimization of Cr content of metastable beta-type Ti-Cr alloys with changeable Young’s modulus for spinal fixation

H. Liu, M. Niinomi, M. Nakai, J. Hieda, K. Cho, Changeable Young’s modulus with large elongation-

ro

[3]

of

applications, Acta Biomater. 8 (2012) 2392–2400. doi:10.1016/j.actbio.2012.02.010.

to-failure in β-type titanium alloys for spinal fixation applications, Scr. Mater. 82 (2014) 29–32.

[4]

-p

doi:10.1016/j.scriptamat.2014.03.014.

H. Liu, M. Niinomi, M. Nakai, K. Cho, K. Narita, M. Sen, H. Shiku, T. Matsue, Mechanical properties

re

and cytocompatibility of oxygen-modified β-type Ti-Cr alloys for spinal fixation devices, Acta

[5]

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Biomater. 12 (2015) 352–361. doi:10.1016/j.actbio.2014.10.014. H. Liu, M. Niinomi, M. Nakai, K. Cho, Athermal and deformation-induced ω-phase transformations in biomedical beta-type alloy Ti-9Cr-0.2O, Acta Mater. 106 (2016) 162–170.

[6]

ur na

doi:10.1016/j.actamat.2016.01.008.

W.F. Ho, T.Y. Chiang, S.C. Wu, H.C. Hsu, Mechanical properties and deformation behavior of cast binary Ti-Cr alloys, J. Alloys Compd. 468 (2009) 533–538. doi:10.1016/j.jallcom.2008.01.046. H.C. Hsu, S.C. Wu, T.Y. Chiang, W.F. Ho, Structure and grindability of dental Ti-Cr alloys, J. Alloys

Jo

[7]

Compd. 476 (2009) 817–825. doi:10.1016/j.jallcom.2008.09.116.

[8]

W.F. Ho, T.Y. Chiang, S.C. Wu, H.C. Hsu, Evaluation of low-fusing porcelain bonded to dental cast Ti-Cr alloys, J. Alloys Compd. 474 (2009) 505–509. doi:10.1016/j.jallcom.2008.06.152.

[9]

K.C. Chen, S.M. Allen, J.D. Livingston, Factors Affecting the Room-Temperature Mechanical Properties of TiCr2-Base Laves Phase Alloys, Mater. Sci. Eng. A. 242 (1998) 162–173.

25

doi:10.1016/S0921-5093(97)00526-1. [10]

R. Banerjee, P.C. Collins, H.L. Fraser, Phase Evolution in Laser-Deposited Titanium-Chromium Alloys, Metall. Mater. Trans. A. 33 (2002) 2129–2138. doi:10.1007/s11661-002-0044-2.

[11]

S. Amira, S.F. Santos, J. Huot, Hydrogen sorption properties of Ti-Cr alloys synthesized by ball milling and cold rolling, Intermetallics. 18 (2010) 140–144. doi:10.1016/j.intermet.2009.07.004.

[12]

P. Frost, W. Parris, L. Hirsch, J. Doig, C. Schwartz, Isothermal transformation of titanium-

K.C. Chen, S.M. Allen, J.D. Livingston, Microstructures of two-phase Ti–Cr alloys containing the

ro

[13]

of

chromium alloys, Trans. Am. Soc. Met. 46 (1954) 1056–1074.

TiCr2 Laves phase intermetallic, J. Mater. Res. 12.6 (1997) 1472–1480.

[14]

-p

doi:10.1557/JMR.1997.0203.

Y.Z. Zhang, C. Meacock, R. Vilar, Laser powder micro-deposition of compositional gradient Ti-Cr

R. Fleischer, R. Zabala, Mechanical properties of Ti-Cr-Nb alloys and prospects for high-

lP

[15]

re

alloy, Mater. Des. 31 (2010) 3891–3895. doi:10.1016/j.matdes.2010.02.052.

temperature applications, Metall. Mater. Trans. A. 21 (1990) 2149–2154. doi:10.1007/BF02647875.

M. Abdel-Hady, K. Hinoshita, M. Morinaga, General approach to phase stability and elastic

ur na

[16]

properties of β-type Ti-alloys using electronic parameters, Scr. Mater. 55 (2006) 477–480. doi:10.1016/j.scriptamat.2006.04.022. H. Liu, M. Niinomi, M. Nakai, K. Cho, β-Type titanium alloys for spinal fixation surgery with high

Jo

[17]

Young’s modulus variability and good mechanical properties, Acta Biomater. 24 (2015) 361–369. doi:10.1016/j.actbio.2015.06.022.

[18]

H. Sahasrabudhe, A. Bandyopadhyay, In situ reactive multi-material Ti6Al4V-calcium phosphatenitride coatings for bio-tribological applications, J. Mech. Behav. Biomed. Mater. 85 (2018) 1–11. doi:10.1016/J.JMBBM.2018.05.020.

26

[19]

K. Stenberg, S. Dittrick, S. Bose, A. Bandyopadhyay, Influence of simultaneous addition of carbon nanotubes and calcium phosphate on wear resistance of 3D-printed Ti6Al4V, J. Mater. Res. 33 (2018) 2077–2086. doi:10.1557/jmr.2018.234.

[20]

S.F. Robertson, A. Bandyopadhyay, S. Bose, Titania nanotube interface to increase adhesion strength of hydroxyapatite sol-gel coatings on Ti-6Al-4V for orthopedic applications, Surf. Coatings Technol. 372 (2019) 140–147. doi:10.1016/J.SURFCOAT.2019.04.071. L.D. Bobbio, B. Bocklund, A. Reichardt, R. Otis, J.P. Borgonia, R.P. Dillon, A.A. Shapiro, B.W.

of

[21]

ro

McEnerney, P. Hosemann, Z.-K. Liu, A.M. Beese, Analysis of formation and growth of the σ phase in additively manufactured functionally graded materials, J. Alloys Compd. (2019) 151729.

[22]

-p

doi:10.1016/J.JALLCOM.2019.151729.

D.R. Waryoba, J.S. Keist, C. Ranger, T.A. Palmer, Microtexture in additively manufactured Ti-6Al-

lP

doi:10.1016/J.MSEA.2018.07.098.

re

4V fabricated using directed energy deposition, Mater. Sci. Eng. A. 734 (2018) 149–163.

[23]

D. Gu, Laser additive manufacturing of high-performance materials, Springer, 2015.

[24]

K.I. Schwendner, R. Banerjee, P.C. Collins, C.A. Brice, H.L. Fraser, Direct laser deposition of alloys

ur na

from elemental powder blends, Scr. Mater. 45 (2001) 1123–1129. doi:10.1016/S13596462(01)01107-1. [25]

R. Banerjee, P.C. Collins, D. Bhattacharyya, S. Banerjee, H.L. Fraser, Microstructural evolution in

Jo

laser deposited compositionally graded α/β titanium-vanadium alloys, Acta Mater. 51 (2003) 3277–3292. doi:10.1016/S1359-6454(03)00158-7.

[26]

P.C. Collins, R. Banerjee, S. Banerjee, H.L. Fraser, Laser deposition of compositionally graded titanium–vanadium and titanium–molybdenum alloys, Mater. Sci. Eng. A. 352 (2003) 118–128. doi:10.1016/S0921-5093(02)00909-7.

[27]

R. Banerjee, S. Nag, H.L. Fraser, A novel combinatorial approach to the development of beta

27

titanium alloys for orthopaedic implants, Mater. Sci. Eng. C. 25 (2005) 282–289. doi:10.1016/j.msec.2004.12.010. [28]

D. Thomas, The Development of Design Rules for Selective Laser Melting, University of Wales Institute, Cardiff., 2009.

[29]

P. Regenfuss, R. Ebert, H. Exner, Laser Micro Sintering – a Versatile Instrument for the Generation of Microparts, Laser Tech. J. 4 (2007) 26–31. doi:10.1002/latj.200790139. S. Yang, Y.F. Zhao, Additive manufacturing-enabled design theory and methodology: a critical

of

[30]

[31]

ro

review, Int. J. Adv. Manuf. Technol. 80 (2015) 327–342. doi:10.1007/s00170-015-6994-5. S. Nag, R. Banerjee, H.L. Fraser, A novel combinatorial approach for understanding

-p

microstructural evolution and its relationship to mechanical properties in metallic biomaterials, Acta Biomater. 3 (2007) 369–376. doi:10.1016/j.actbio.2006.08.005.

G.H. Narayanan, T.S. Luhman, T.F. Archbold, R. Taggart, D.H. Polonis, A phase separation reaction

re

[32]

0800(71)90063-2. [33]

lP

in a binary titanium-chromium alloy, Metallography. 4 (1971) 343–358. doi:10.1016/0026-

M. Hattori, S. Takemoto, M. Yoshinari, E. Kawada, Y. Oda, Effect of chromium content on

ur na

mechanical properties of casting Ti-Cr alloys, Dent. Mater. J. 29 (2010) 570–4. doi:10.4012/dmj.2009-118. [34]

P. Samimi, Y. Liu, I. Ghamarian, D.A. Brice, P.C. Collins, A new combinatorial approach to assess

Jo

the influence of alloy composition on the oxidation behavior and concurrent oxygen-induced phase transformations for binary Ti-xCr alloys at 650°C, Corros. Sci. 97 (2015) 150–160. doi:10.1016/j.corsci.2015.05.002.

[35]

D.J. Lin, J.H. Chern Lin, C.P. Ju, Effect of omega phase on deformation behavior of Ti-7.5Mo-xFe alloys, Mater. Chem. Phys. 76 (2002) 191–197. doi:10.1016/S0254-0584(01)00511-9.

[36]

S. Guo, Q. Meng, G. Liao, L. Hu, X. Zhao, Microstructural evolution and mechanical behavior of

28

metastable β-type Ti–25Nb–2Mo–4Sn alloy with high strength and low modulus, Prog. Nat. Sci. Mater. Int. 23 (2013) 174–182. doi:10.1016/j.pnsc.2013.03.008. [37]

S. Dubinskiy, A. Korotitskiy, S. Prokoshkin, V. Brailovski, In situ X-ray diffraction study of athermal and isothermal omega-phase crystal lattice in Ti-Nb-based shape memory alloys, Mater. Lett. 168 (2016) 155–157. doi:10.1016/j.matlet.2016.01.012.

[38]

W.F. Ho, Effect of omega phase on mechanical properties of Ti-Mo alloys for biomedical

B. Zhang, J. Wang, W. Xiaojing, C. Weijie, A study on the beta and omega phases in a Ti-Al-Cr

ro

[39]

of

applications, J. Med. Biol. Eng. 28 (2008) 47–51.

alloy, Scr. Metall. Mater. 30 (1994) 399–404. doi:10.1016/0956-716X(94)90593-2. S. Nag, R. Banerjee, R. Srinivasan, J.Y. Hwang, M. Harper, H.L. Fraser, ω-Assisted nucleation and

-p

[40]

growth of α precipitates in the Ti-5Al-5Mo-5V-3Cr-0.5Fe β titanium alloy, Acta Mater. 57 (2009)

K. Madangopal, J.B. Singh, S. Banerjee, The nature of self-accommodation in Ni-Ti shape memory

lP

[41]

re

2136–2147. doi:10.1016/j.actamat.2009.01.007.

alloys, Scr. Metall. Mater. 29 (1993) 725–728. doi:10.1016/0956-716X(93)90215-E. [42]

D. Srivastava, K. Madangopal, S. Banerjee, S. Ranganathan, Self accomodation morphology of

ur na

martensite variants in Zr-2.5 wt% Nb alloy, Acta Metall. Mater. 41 (1993) 3445–3454. doi:10.1016/0956-7151(93)90224-G. G. Lütjering, J.C. Williams, Titanium, Springer Science & Business Media, Berlin, 2007.

[44]

G. Welsch, R. Boyer, E. Collings, Materials properties handbook: titanium alloys, 1993.

Jo

[43]

[45]

S. Banerjee, P. Mukhopadhyay, Phase transformations: examples from titanium and zirconium alloys, 2010th ed., Elsevier, 2010.

[46]

A. Goldenstein, A. Metcalfe, W. Rostoker., The effect of stress on the eutectoid decomposition of titanium–chromium alloys, in: Trans. ASM 51, 1959: pp. 1036–1053.

[47]

B. Zheng, Y. Zhou, J.E. Smugeresky, J.M. Schoenung, E.J. Lavernia, Thermal behavior and

29

microstructural evolution during laser deposition with laser-engineered net shaping: Part I. Numerical calculations, Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 39 (2008) 2228–2236. doi:10.1007/s11661-008-9557-7. [48]

B. Zheng, Y. Zhou, J.E. Smugeresky, J.M. Schoenung, E.J. Lavernia, Thermal behavior and microstructure evolution during laser deposition with laser-engineered net shaping: part II. Experimental investigation and discussion, Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 39

P. Michaleris, Modeling metal deposition in heat transfer analyses of additive manufacturing

ro

[49]

of

(2008) 2237–2245. doi:10.1007/s11661-008-9566-6.

processes, Elsevier. 86 (2014) 51–60. doi:10.1016/j.finel.2014.04.003.

S.M. Kelly, S.L. Kampe, Microstructural evolution in laser-deposited multilayer Ti-6Al-4V builds:

-p

[50]

Part II. Thermal modeling, Metall. Mater. Trans. A. 35 (2004) 1869–1879. doi:10.1007/s11661-

S.M. Kelly, S.L. Kampe, Microstructural evolution in laser-deposited multilayer Ti-6Al-4V builds:

lP

[51]

re

004-0095-7.

Part I. Microstructural characterization, Metall. Mater. Trans. A. 36 (2004) 1861–1867. doi:10.1007/s11661-004-0094-8.

J. Heigel, P. Michaleris, E.R. Manufacturing, Thermo-mechanical model development and

ur na

[52]

validation of directed energy deposition additive manufacturing of Ti–6Al–4V, Addit. Manuf. 5 (2015) 9–19. doi:10.1016/j.addma.2014.10.003. L. Costa, R. Vilar, T. Reti, A.. Deus, Rapid tooling by laser powder deposition: process simulation

Jo

[53]

using finite element analysis, Acta Mater. 53 (2005) 3987–3999. doi:10.1016/j.actamat.2005.05.003.

[54]

P. Peyre, P. Aubry, R. Fabbro, R. Neveu, A. Longuet, Analytical and numerical modelling of the direct metal deposition laser process, J. Phys. D. Appl. Phys. 41 (2008) 025403. doi:10.1088/00223727/41/2/025403.

30

[55]

L. Wang, S. Felicelli, Analysis of thermal phenomena in LENSTM deposition, Mater. Sci. Eng. A. 435 (2006) 625–631. doi:10.1016/j.msea.2006.07.087.

[56]

A.D. Niessen, F.R. De Boer, R.D. Boom, P.F. De Chatel, Model predictions for the enthalpy of formation of transition metal alloys II, Calphad. 7 (1983) 51–70. doi:10.1016/03645916(83)90030-5.

[57]

W. Hofmeister, M. Griffith, U. 2001, Solidification in direct metal deposition by LENS processing,

R. Dong, J. Li, J. Fan, H. Kou, B. Tang, Precipitation of α phase and its morphological evolution

ro

[58]

of

JOM. 53 (2001) 30–34. doi:10.1007/s11837-001-0066-z.

during continuous heating in a near β titanium alloy Ti-7333, Mater. Charact. 132 (2017) 199–

[59]

-p

204. doi:10.1016/j.matchar.2017.07.032.

P.L. Narayana, S. Lee, S.W. Choi, C.L. Li, C.H. Park, J.T. Yeom, N.S. Reddy, J.K. Hong,

re

Microstructural response of β-stabilized Ti–6Al–4V manufactured by direct energy deposition, J.

[60]

lP

Alloys Compd. 811 (2019). doi:10.1016/j.jallcom.2019.152021. S. Nag, Y. Zheng, R.E.A. Williams, A. Devaraj, A. Boyne, Y. Wang, P.C. Collins, G.B. Viswanathan, J.S. Tiley, B.C. Muddle, R. Banerjee, H.L. Fraser, Non-classical homogeneous precipitation

ur na

mediated by compositional fluctuations in titanium alloys, Acta Mater. 60 (2012) 6247–6256. doi:10.1016/j.actamat.2012.07.033. [61]

F. Prima, P. Vermaut, G. Texier, D. Ansel, T. Gloriant, Evidence of α-nanophase heterogeneous

Jo

nucleation from ω particles in a β-metastable Ti-based alloy by high-resolution electron microscopy, Scr. Mater. 54 (2006) 645–648. doi:10.1016/j.scriptamat.2005.10.024.

[62]

M. Ahmed, T. Li, G. Casillas, J.M. Cairney, D. Wexler, E. V. Pereloma, The evolution of microstructure and mechanical properties of Ti-5Al-5Mo-5V-2Cr-1Fe during ageing, J. Alloys Compd. 629 (2015) 260–273. doi:10.1016/j.jallcom.2015.01.005.

[63]

Y. Ikematsu, M. Doi, T. Miyazaki, Phase decomposition of liquid-quenched β-type Ti-Cr alloys, J.

31

Mater. Sci. 26 (1991) 2071–2075. doi:10.1007/BF00549169. [64]

T. Luhman, R. Taggart, D. Polonis, Correlation of superconducting properties with the beta to omega phase transformation in Ti-Cr alloys, Scr. Metall. 3 (1969) 777–782. doi:10.1016/00369748(69)90178-1.

[65]

S.L. Semiatin, S.L. Knisley, P.N. Fagin, D.R. Barker, F. Zhang, Microstructure evolution during alpha-beta heat treatment of Ti-6Al-4V, Metall. Mater. Trans. A. 34 (2003) 2377–2386.

C. Qiu, Q. Liu, Multi-scale microstructural development and mechanical properties of a selectively

ro

[66]

of

doi:10.1007/s11661-003-0300-0.

Jo

ur na

lP

re

-p

laser melted beta titanium alloy, Addit. Manuf. 30 (2019). doi:10.1016/j.addma.2019.100893.

32